High-performance CuFe₂O₄ epitaxial thin films with enhanced ferromagnetic resonance properties

Abstract

Highly epitaxial thin films of copper ferrite (CuFe₂O₄) have been fabricated on MgAl₂O₄ (001) substrates at a growth temperature of 400 °C for the first time, we believe, through a radio-frequency sputtering method. Structural analyses through high-resolution X-ray diffraction (HRXRD), Raman spectroscopy, and high-resolution transmission electron microscopy (HRTEM) all confirm the tetragonal spinel phase of CuFe₂O₄ epitaxial film and high tetragonal distortion (c/a = 1.08) of its unit cells. The 50 nm-thick T-CuFe₂O₄ epitaxial film shows unique soft magnetism with small coercivity of 23 Oe and decreased magnetization. However, a superior ferromagnetic linewidth of only ∼93 Oe in the post-annealed T-CuFe₂O₄ film compared with a linewidth of ∼1500 Oe in T-CuFe₂O₄ single crystal bulk material is also observed, which indicates that epitaxial growth of oxide thin films combined with proper heat-treatment-induced cation engineering can impose novel functionalities.

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1. Introduction

Multifunctional oxides containing transition metal elements have attracted considerable attention for fundamental and application research owing to their novel applications over a wide area.^1–4 Spinel ferrites have been extensively investigated for their unique physical properties in magnetic, electric and optic aspects.^5–7 High Curie temperature, good chemical stability, and high catalytic activity make spinel ferrites potential candidates for applications in spintronics, heterogeneous catalysts, and gas sensors.^8–11 Usually, the vast majority of the spinel ferrite crystals (MFe₂O₄, M = Co, Ni, Zn, Mg, etc.) have a cubic symmetry, which is very stable and resistant to temperature. However, copper ferrite (CuFe₂O₄) is very special and shows two symmetries, cubic (space group: Fd3̄m) and tetragonal (space group: I4₁/amd). The structure transformation from high symmetry (Fd3̄m) to low symmetry (I4₁/amd) induced by Jahn–Teller distortion has been extensively studied for decades.^12,13 Appropriate heat-treatment can stabilize both the cubic and tetragonal copper ferrite (C-CFO and T-CFO) at room temperature. According to previous studies, the structural change will dramatically alter its electric, magnetic and photonic properties.^14–16 Aside from the above intriguing structural and physical properties, copper ferrite also exhibits efficiency in catalytic removal of nitrogen oxides and diesel particulates, which helps build a sustainable environment.¹⁷ Therefore, the versatility of copper ferrite has been studied intensively in the form of ceramics, nanoparticles, and polycrystalline films.^18–20 Only a few works have reported the CFO epitaxial thick films grown by the liquid-phase epitaxy (LPE) method.^21,22 However, CFO epitaxial thin films fabricated by modern film growth technologies have not been reported and studied so far.²³ It is well known that epitaxial thin films with least lattice defects, large epitaxial strain, and obvious interface effect will have novel properties, which are hardly seen in other forms of material.^24,25 Therefore, the studies on the epitaxial growth of CFO film and its related properties are of great interest in our research.

In this work, we report on what we believe to be the first epitaxial growth of CFO thin films with excellent quality on MgAl₂O₄ (MAO) (001) substrate using a radio frequency (RF) sputtering system. The 50 nm-thick CFO/MAO (001) film with high tetragonality of c/a = 1.08 shows many novel properties, which are closely related to the process of film preparation. The soft magnetism of the CFO film is demonstrated by magnetic hysteresis loops and ferromagnetic resonance measurement, which deviate from properties of typical T-CFO bulk material and polycrystalline film. The peak-to-peak linewidth of ∼93 Oe under 9.2 GHz of CFO film also shows much improvement compared with the linewidth of ∼1500 Oe in CFO single crystal. It is indicated that ferrite thin films with careful after-deposition heat-treatment may have better performance than even single crystal bulk material, which may pave an effective way to novel functionalities for oxides.

2. Experimental

2.1 Synthesis of ferrite ceramic target

An RF sputtering system was employed to grow the CuFe₂O₄ (CFO) thin films. The stoichiometric CFO ceramic target was synthesized by a conventional solid reaction method with initial reactants CuO (99.99%, provided by Sinopharm Chemical Reagent Co., Ltd) and Fe₂O₃ (99.99%, provided by MTI KJ Group) powders (CuO : Fe₂O₃ = 1 : 2). X-ray diffraction (XRD) measurement showed that the prepared CFO target was pure tetragonal spinel phase with lattice parameters a = 8.239 Å and c = 8.658 Å. The c/a ratio of the T-CFO target is 1.05, slightly smaller than the reported tetragonality of 1.06 for the T-CFO bulk material (a = 8.202 Å and c = 8.730 Å). This discrepancy in lattice parameters can be understood by the lower sintering temperature of 900 °C than used in the literature, because a lower sintering temperature will lead to a lower inversion parameter in ceramic CFO, which results in more reduced Jahn–Teller distortion and smaller tetragonality.

2.2 Fabrication of thin films

Before the thin film growth process, the CFO target was pre-sputtered for over 12 hours to eliminate potential contaminant on the surface. The single-crystal MgAl₂O₄ (MAO) (001) substrate with lattice parameter of 8.083 Å was selected to stabilize T-CFO film owing to the small lattice mismatch of 1.45%. The growth conditions were set at a growth temperature of 400 °C under a working pressure of 0.5 mbar with a mixed ambient atmosphere of Ar and O₂ in a ratio of 1 : 1 during the deposition process. The thickness was estimated to be about 50 nm with a growth rate of 2.74 nm h⁻¹ and RF power of 100 W. After finishing the growth, the CFO thin films were annealed at 400 °C for 10 min in 400 mbar Ar/O₂ atmosphere to eliminate oxygen vacancies in the films. After that the heater was turned off and the sample was rapidly cooled down to room temperature in 20 minutes. To further optimize ferromagnetic resonance properties and lower magnetic inhomogeneity, the as-grown CFO films (AG-CFO sample) were further annealed in a furnace at 400 °C for 15 minutes. After that the films were cooled slowly at a rate of ∼2 °C min⁻¹ to room temperature (CFO film that underwent this process was denoted SC-CFO sample).

2.3 Characterization

The epitaxial nature, crystalline quality, surface morphology, and strain state of the CFO films were characterized by the high-resolution X-ray diffraction system PANalytical X'pert MRD with a wavelength of 1.540598 Å. To accurately determine the thickness of the epitaxial thin film, an X-ray reflection (XRR) curve ranging from 0.1° to 2.0° was recorded by using the ω–2θ scanning method with 1/8 receiving slit. Prior to XRR measurement, the X-ray reflection peak was searched by setting a small detection angle 2θ of 0.8° and performing an ω scan. The atomic force microscopy (AFM) image of 1 μm × 1 μm area was scanned by using a Bruker Dimension Icon. The short-range ion relationships were measured by Horiba hr800 Raman spectroscopy using 633 nm excitation sources. Bright field (BF) image and selected-area electron diffraction were recorded on a JEOL 2100 transmission electron microscope, and a scanning transmission electron microscopy (STEM) investigation was performed on a JEOL ARM200F microscope. The room temperature magnetic properties of M–H hysteresis loops and ferromagnetic resonance response were performed using Quantum Design VSM and a JES-FA 200 EPR system with 9.2 GHz microwave source, respectively.

3. Results and discussion

The structural properties (crystalline quality, crystallographic relationship, thickness, and strain relaxation states) of the sputtering-grown CFO film on MAO substrate were studied by XRD θ–2θ scan, rocking curve, ϕ scan, reflectivity, and reciprocal space mapping (RSM). Fig. 1(a) shows a typical θ–2θ scanning pattern of SC-CFO sample grown on the (001) MAO substrate. Only the CFO (004) peak and the MAO (004) peak can be found, demonstrating that the CFO film is c-axis oriented without any secondary phases. The clear satellite peaks around the CFO (004) reflection shown in the inset of Fig. 1(a) indicate smooth film surface morphology and sharp film–substrate interface. The film thickness was evaluated by using X-ray reflectivity (XRR) measurement, as shown in Fig. 1(b). The wide-angle range interference pattern of the XRR curve demonstrates very smooth surface morphology. The film thickness of ∼50 nm was evaluated from the formula T = nλ/(θ_m+n − θ_m), where T, λ, θ_m, and n are the film thickness, the wavelength of the incident X-rays, the incident angle of the mth interference peak, and the differential of the interference orders between peaks, respectively. The surface roughness value (root mean square) measured by using atomic force microscopy (seen in the inset of Fig. 1(b)) is only 0.144 nm, showing atomic-flat surface morphology of the CFO film. The rocking curve measurement from the CFO (004) reflection shows that the full width at half maximum (FWHM) is only about 0.1°, as shown in Fig. 1(c), revealing the excellent crystalline quality of the CFO film. For better understanding of the in-plane crystallographic relationship between the CFO film and the MAO substrate, ϕ scans were performed, as shown in Fig. 1(d). The four-fold symmetry and sharp peaks in the ϕ scans of the CFO films further indicate that the CFO film has good crystalline quality and epitaxial nature. The orientation relationship between the CFO film and the MAO substrate can be determined as [100]_CFO//[100]_MAO and (001)_CFO//(001)_MAO.

Fig. 1 High resolution XRD investigations on 50 nm-thick CFO/MAO (001) film. (a) A wide-angle θ–2θ scan. The detailed θ–2θ scan in the inset shows clear satellite peaks around CFO (004) diffraction. (b) Normalized X-ray reflectivity. The inset is an AFM image of CFO film showing smooth surface morphology. (c) Rocking curve recorded around the CFO (004) diffractions with FWHM = 0.1°. (d) ϕ Scans taken around the {101} diffractions from CFO film and MAO substrate.

The reciprocal space mapping (RSM) technique was used to obtain more details of the structure, such as lattice parameters and strain state. RSM patterns were obtained around the asymmetric (206) diffractions of the SC-CFO film and MAO substrate as shown in Fig. 2(a). Similar to the θ–2θ scan, clear interference patterns symmetrically distributed around the sharp primary peak of CFO (206) reflection are also seen. The positions of CFO (206) RSM patterns lie exactly at the horizontal axis across MAO (206) RSM, indicating the coherency of the in-plane lattice constants (8.083 Å) of film and substrate. The large distance between CFO and MAO (206) peak positions suggests a much larger out-of-plane lattice constant (8.754 Å) of CFO film and a high c/a ratio of 1.083. This large tetragonality surpasses the highest c/a ratio of 1.06 in tetragonal CFO bulk material showing in-plane compressive strain induced by substrate.¹⁹ To further confirm the structural features of the SC-CFO film and investigate its short-range structure, intrinsic Raman spectroscopy analysis was employed. In order to obtain the real Raman information of the CFO layer, the signals coming from MAO substrate have already been subtracted from the whole spectrum of CFO/MAO as illustrated in Fig. 2(b). According to factor group analysis, 10 Raman active modes exist in the tetragonal CFO, which can be represented by 2T(B_1g ⊕ E_g) + L(E_g) + ν₁(A_1g) + 2ν₂(A_1g ⊕ B_2g) + 2ν₃(B_1g ⊕ E_g) + 2ν₄(B_1g ⊕ E_g).²⁶ The tetragonal CFO has more active modes than the typical cubic phase spinel ferrites with only 5 active modes. The 8 visible Raman modes in the CFO film are broad and overlap with each other, which is similar to previous results.^26,27 To accurately determine the Raman shift of each visible mode, a Gaussian peak function was used to fit multiple Raman peaks of CFO film. Fig. 2(b) suggests the fitting curve matches well with the experimental curve so that all the visible modes can be identified and given special assignments. Therefore, the tetragonal spinel structure of CFO film can be further confirmed. Moreover, the obvious blue shift of the A_1g mode (727 cm⁻¹) of the T-CFO film compared with the A_1g mode frequency (691 cm⁻¹) of the T-CFO bulk material reveals that the Fe(Cu)–O bond length in AO₄ tetrahedron shrinks due to epitaxial strain from the substrate, which is in agreement with our RSM data.

Fig. 2 (a) Reciprocal space mapping around the (206) diffractions of the CFO film and the MAO substrate. (b) Intrinsic Raman spectrum of CFO layer after subtracting the fluorescent signal of MAO substrate showing short-range structure information of the CFO film. A simulated line obtained by employing Gaussian peak function is used to fit the experimental line and determine the exact shift of each Raman peak.

Additionally, the epitaxial behavior and microstructural properties of SC-CFO film on (001) MAO were studied by means of (scanning) transmission electron microscopy (TEM). Fig. 3(a) shows a typical low-magnification bright field image of the CFO film on the MAO substrate. The film thickness of ∼50 nm, sharp CFO–MAO interface (arrowed), and smooth surface morphology are consistent with the results from XRR and AFM measurement. Fig. 3(b) shows the selected-area electron diffraction (SAED) patterns from the area covering the CFO film and the MAO substrate. The large splitting of diffraction peaks from CFO and MAO is clearly visible in Fig. 3(b), as denoted by a horizontal white arrow. According to the SAED patterns, the epitaxial orientation relationship of [100]_CFO//[100]_MAO and (001)_CFO//(001)_MAO can again be determined between CFO and MAO. Fig. 3(c) is an atomic-resolution high-angle annular dark field (HAADF) image of the CFO/MAO interface region viewed along the [100] of MAO. The interdiffusion between copper-based oxides and substrate is not observed because the very low epitaxial growth temperature of 400 °C was adopted during the film growth process. Lattice-misfit-induced dislocations, as well as other kinds of structure defects, can hardly be observed either, further demonstrating the perfect crystalline quality of our CFO film. In short, the CFO films can be epitaxially grown by sputtering at low temperature with few defects and perfect crystalline quality.

Fig. 3 (a) A low-magnification bright field (BF) image of the CFO/MAO system. (b) SAED pattern from the area including the CFO film and the MAO substrate. (c) An atomic-resolution HAADF image of the CFO/MAO interface region viewed along the zone axis [100] of MAO.

In order to analyze the correlation between the microstructure and physical properties and the effect of post-deposition heat-treatment on physical properties, the magnetic properties of the epitaxial CFO films on (001) MAO substrates were measured. Fig. 4(a) shows the magnetic hysteresis loops measured along the in-plane direction of the CFO film at 300 K. The T-CFO film exhibits well-defined soft magnetism with smaller coercivity of only 23 Oe in SC-CFO film according to the enlarged loops in the inset of Fig. 4(a). This coercivity value is much smaller than those of typical T-CFO bulk material and polycrystalline films in the literature, which exhibit large coercivity (typically around several hundreds of oersteds) caused by comparatively large uniaxial anisotropy.^28,29 Dozens of T-CFO/MAO (001) epitaxial films with different thickness grown by us at different deposition pressure and substrate temperature all show similar small coercivities, suggesting the soft magnetism measured is unrelated to particular growth conditions. A possible explanation to this anomaly may lie in the highly epitaxial structure and large tetragonal distortion in our CFO film. Moreover, owing to the demagnetization field and compressive-strain-induced in-plane magnetic anisotropy, CFO/MAO (001) thin film with negative magnetostriction does not show perpendicular magnetic anisotropy. The SC-CFO film shows 12% enhancement of saturation magnetization (M_s) compared with that of AG-CFO film. However, the value of M_s (50 emu cm⁻³) measured in SC-CFO is much decreased compared with that of T-CFO bulk material (135 emu cm⁻³ or 1700 Gs).³⁰ Several factors could cause the decreased magnetization for typical spinel ferrite films, such as oxygen vacancies, structure defects (antiphase boundaries, grain boundaries, dislocations, etc.) and inversion parameters. As we used a pretty high deposition pressure of 0.5 mbar and post-annealing treatment in air to process the CFO film, the decreased magnetization is unlikely to be caused by oxygen vacancies. As structure imperfection is demonstrated to be minimal it should not be the dominant factor causing the magnetization reduction. The inversion parameter is particularly important for the measured magnetization in mixed spinel like copper ferrite, which can be denoted as [Cu_1−λFe_λ]_Td{Cu_λFe_2−λ}_OhO₄ (λ is the inversion parameter).^14,30 The overall molecular magnetization in copper ferrite can be described by the equation M_s = M_Oh − M_Td = 9 − 8λ μ_B (noting the Cu²⁺ and Fe³⁺ ions contribute to molecular moments of 1 μ_B and 5 μ_B, respectively). The more inverse spinel structure of copper ferrite is present, the less magnetization can be measured. A. Yang et al. have studied the effect of deposition pressure on magnetization and inversion parameter in CFO film, and they found films grown above 90 mTorr had increased inversion parameter and decreased magnetization.²³ We adopted a much higher deposition pressure, which should also induce a high inversion parameter in our film. Moreover, the fact that coherent growth mode was retained up to the thickness of 50 nm in the CFO film without relaxation also favors the assumption of high inversion parameter, since the lattice misfit between film and substrate will become smaller when the unit cell of CFO film with higher inversion parameter experiences more severe Jahn–Teller distortion. The reason that SC-CFO film shows enhancement of M_s compared with that of AG-CFO film can also be understood by cation redistribution induced by heat-treatment. As indicated by Fig. 4(b), post-deposition heat-treatment does not change strain states or crystalline quality, since both θ–2θ scan and rocking curve present the same result in both AG-CFO and SC-CFO films. However, previous study in CFO shows that heat-treatment is more effective in changing M_s by changing its cation distribution.¹⁴ Thus it is reasonable to believe the change of M_s mainly comes from cation redistribution in the CFO lattice. It is worth noting that direct growth of CFO films at higher temperature (above 400 °C) without further annealing also causes slight enhancement of both magnetization and coercivity as a result of growth-temperature-induced cation redistribution, which is in accordance with previous reports in the literature.²³

Fig. 4 (a) Magnetic hysteresis loops of the 50 nm-thick CFO films measured along the in-plane direction at 300 K. The inset shows the small coercivities in an enlarged area of the M–H loops. (b) Through θ–2θ scan and rocking curve, no visible shift of diffraction angle and change of crystalline quality are observed in CFO film after post-deposition heat-treatment. (c) Ferromagnetic resonance response of the CFO films with the magnetic field applied in the film plane. The black line shows a single and strong FMR absorption in SC-CFO film while the blue line shows multiple and weak FMR absorptions in AG-CFO film with 250-fold magnification of data.

To evaluate the magnetization dynamic and magnetic homogeneity of CFO films, room temperature ferromagnetic resonance measurement at 9.2 GHz with in-plane configuration was performed. The black line in Fig. 4(c) shows strong ferromagnetic absorption in SC-CFO film. The measured peak-to-peak linewidth is ∼93 Oe, which is much narrower than the value for T-CFO single crystal bulk material and many other spinel ferrite epitaxial films.³¹ The damping factor α of SC-CFO films can be calculated through the equation α = 3ΔHγ/2ω, which is deduced from the Landau–Lifshitz–Gilbert (LLG) equation.³² ΔH, γ and ω stand for the FMR linewidth, gyromagnetic ratio and resonance frequency, respectively. As γ equals 2.8 MHz Oe⁻¹, the calculated damping factor is 0.0245, which is quite small when compared with that of many ternary spinel ferrite thin films.³¹ T. Okamura et al. have measured the ferromagnetic resonance in a T-CFO single crystal. Similar to us, their single crystal sample also exhibited decreased magnetization (only about 300 Gs cm⁻³). The narrowest linewidth they measured in a spherical sample was around 1500 Oe.³³ Typically, ferromagnetic linewidth of ferrites is broadened in epitaxial films and its value in films is 1 to 2 orders larger than that in single crystal.³¹ Our measured linewidth of 93 Oe in T-CFO film shows much improvement compared with their results. This anomaly in CFO prompted us to refer to the literature, and we found that the crystallinity of copper ferrite single crystal suffers from the phenomenon of micro twinning, which is a natural habit when a copper ferrite single crystal slowly cools down and passes through the phase transition temperature (around 660–700 K), induced by the Jahn–Teller effect.^19,34 We speculate that the degenerated crystalline quality in the copper ferrite single crystal has a large influence on the broadening of the linewidth. Moreover, the importance of post-deposition heat-treatment in reducing magnetic inhomogeneity and helping enhance ferromagnetic resonance of CFO film was revealed by testing FMR response of CFO film without post-annealing process (AG-CFO film). This suggests that AG-CFO film shows multiple, weak, and broadened FMR peaks (shown as the blue line of Fig. 4(c)). A previous study had demonstrated that thermal history can easily influence cation distribution in copper ferrite because of the low activation energy (∼0.1 eV) when Cu²⁺ ions change position.¹⁹ And the way cations are distributed in copper ferrite has a large influence on structural, magnetic and electric properties of copper ferrite.^14,35 The magnetization dynamic process of copper ferrite can also be modified if the local exchange field is changed by the different ways that cations are distributed in the copper ferrite lattice.³⁶ Higher temperatures (above 400 °C) were also adopted to anneal CFO films, and we found that FMR properties do not benefit from higher annealing temperature owing to deterioration of crystalline quality and severe interdiffusion at the interface between film and substrate. At the operating temperature of 400 °C, both the crystalline quality and FMR properties of CFO films are optimized. In short, highly epitaxial tetragonal phase CFO film with appropriate post-annealing heat-treatment shows much improved ferromagnetic resonance compared with that of the single crystal, which suggests that epitaxial growth of functional oxide materials combined with cation engineering may impose superior functionalities.

4. Conclusions

In summary, epitaxial thin films of tetragonal spinel phase copper ferrite with perfect crystalline quality have been coherently grown on MAO (001) substrates, we believe for the first time. Owing to a highly inverse spinel structure, CFO film has larger tetragonal distortion (c/a = 1.08) and lower magnetization than the CFO bulk material. However, unlike many T-CFO polycrystalline films, our sputtering-grown CFO epitaxial film shows unique soft magnetism with very small coercivity of 23 Oe. In-plane ferromagnetic resonance testing also indicates that T-CFO epitaxial film after post-annealing has superior linewidth to that of T-CFO single crystal. It is shown that combining epitaxial growth of functional oxides with cation engineering through appropriate after-deposition heat-treatment can impose superior structural and physical properties, which may pave a feasible way to novel functionalities in oxide thin films.

Acknowledgements

The work was supported by the National Natural Science Foundation of China (no. 61471290, 51390472 and 51202185), the SRFDP-RGC Joint Research Project 2013/14 (20130201140002) and National 973 projects of China (no. 2015CB654603). It was also partially supported by the Fundamental Research Funds for the Central University. We thank Mr Ming-Min Zhu and Prof. Ming Liu for the support of ferromagnetic resonance measurements.

References

1. M. Bibes and A. Barthelemy, Nat. Mater., 2008, 7, 425–426.
2. C. Cen, S. Thiel, J. Mannhart and J. Levy, Science, 2009, 323, 1026–1030.
3. S. Ishiwata, D. Wang, T. Saito and M. Takano, Chem. Mater., 2005, 17, 2789–2791.
4. D. Cao, L. Pan, H. Li, J. Li, X. Wang, X. Cheng, Z. Wang, J. Wang and Q. Liu, RSC Adv., 2016, 6, 66795–66802.
5. Y. Suzuki, Annu. Rev. Mater. Res., 2001, 31, 265–289.
6. M. Alexe, M. Ziese, D. Hesse, P. Esquinazi, K. Yamauchi, T. Fukushima, S. Picozzi and U. Gösele, Adv. Mater., 2009, 21, 4452–4455.
7. R. Sharma, S. Bansal and S. Singhal, RSC Adv., 2016, 6, 71676–71691.
8. U. Luders, M. Bibes, K. Bouzehouane, E. Jacquet, J.-P. Contour, S. Fusil, J.-F. Bobo, J. Fontcuberta, A. Barthélémy and A. Fert, Appl. Phys. Lett., 2006, 88, 082505.
9. H. Y. Zhang, S. T. Gao, N. Z. Shang, C. Wang and Z. Wang, RSC Adv., 2014, 4, 31328–31332.
10. R. B. Kamble and V. L. Mathe, Sens. Actuators, B, 2008, 131, 205–209.
11. Z. Li, X. Lai, H. Wang, D. Mao, C. Xing and D. Wang, J. Phys. Chem. C, 2009, 113, 2792–2797.
12. H. J. Levinstein, F. J. Schnettler and E. M. Gyorgy, J. Appl. Phys., 1965, 36, 1163.
13. R. G. Kulkarni and V. U. Patil, J. Mater. Sci., 1980, 15, 2221–2223.
14. M. Desai, S. Prasad, N. Venkataramani, I. Samajdar, A. K. Nigam and R. Krishnan, J. Appl. Phys., 2002, 91, 2220–2227.
15. X. Zuo, A. Yang, C. Vittoria and V. G. Harris, J. Appl. Phys., 2006, 99, 08M909.
16. M. Kučera, V. Kolinský, S. Višňovský, D. Chvostová, N. Venkataramani, S. Prasad, P. D. Kulkarni and R. Krishnan, J. Magn. Magn. Mater., 2007, 316, e688–e691.
17. Z. Jiang, W. H. Zhang, W. F. Shangguan, X. J. Wu and Y. Teraokag, J. Phys. Chem. C, 2011, 115, 13035–13040.
18. B. K. Chatterjee, K. Bhattacharjee, A. Dey, C. K. Ghosh and K. K. Chattopadhyay, Dalton Trans., 2014, 43, 7930–7944.
19. A. M. Balagurov, I. A. Bobrikov, V. Y. Pomjakushin, D. V. Sheptyakov and V. Y. Yushankhai, J. Magn. Magn. Mater., 2015, 374, 591–599.
20. M. Epifani, J. Arbiol, T. Andreu and J. R. Morante, Cryst. Growth Des., 2010, 10, 5176–5181.
21. D. A. Herman, R. L. White, R. S. Feigelson, B. L. Mattes and H. W. Swarts, AIP Conf. Proc., 1975, 24, 580–581.
22. P. J. M. Van der Straten and R. Metselaar, J. Appl. Phys., 1980, 51, 3236–3240.
23. A. Yang, X. Zuo, L. Chen, Z. Chen, C. Vittoria and V. G. Harris, J. Appl. Phys., 2005, 97, 10G107.
24. M. Liu, C. Ma, G. Collins, J. Liu, C. Chen, C. Dai, Y. Lin, L. Shui, F. Xiang, H. Wang, J. He, J. Jiang, E. I. Meletis and M. W. Cole, ACS Appl. Mater. Interfaces, 2012, 4, 5761–5765.
25. M. Liu, Q. Zou, C. Ma, G. Collins, S. B. Mi, C. L. Jia, H. Guo, H. Gao and C. Chen, ACS Appl. Mater. Interfaces, 2014, 6, 8526–8530.
26. M. D. P. Silva, F. C. Silva, F. S. M. Sinfrônio, A. R. Paschoal, E. N. Silva and C. W. A. Paschoal, J. Alloys Compd., 2014, 584, 573–580.
27. K. Verma, A. Kumar and D. Varshney, Curr. Appl. Phys., 2013, 13, 467–473.
28. S. J. Stewart, R. C. Mercader, R. E. Vandenberghe, G. Cernicchiaro and R. B. Scorzelli, J. Appl. Phys., 2005, 97, 054304.
29. M. Ahmad, M. Desai and R. Khatirkar, J. Appl. Phys., 2008, 103, 013903.
30. A. Yang, Z. Chen, X. Zuo, D. Arena, J. Kirkland, C. Vittoria and V. G. Harris, Appl. Phys. Lett., 2005, 86, 252510.
31. Ü. Özgür, Y. Alivov and H. Morkoç, J. Mater. Sci.: Mater. Electron., 2009, 20, 789–834.
32. K. Zakeri, J. Lindner, I. Barsukov, R. Meckenstock, M. Farle, U. von Hörsten, H. Wende, W. Keune, J. Rocker, S. S. Kalarickal, K. Lenz, W. Kuch, K. Baberschke and Z. Frait, Phys. Rev. B, 2007, 76, 104416.
33. T. Okamura and Y. Kojima, Phys. Rev., 1952, 86, 1040–1041.
34. H. Nagata, T. Miyadai and S. Miyahara, IEEE Trans. Magn., 1972, 8, 451–453.
35. V. V. Parfenov and R. A. Nazipov, Inorg. Mater., 2002, 38, 78–82.
36. K. B. Modi, P. Y. Raval, S. V. Dulera, C. R. Kathad, S. J. Shah, U. N. Trivedi and U. Chandra, AIP Conf. Proc., 2015, 1665, 130011.

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